Method of controlling and refining final grain size in supersolvus heat treated nickel-base superalloys

ABSTRACT

A method of forming a component from a gamma prime precipitation-strengthened nickel-base superalloy. The method entails formulating the superalloy to have a sufficiently high carbon content and forging the superalloy at sufficiently high local strain rates so that, following a supersolvus heat treatment, the component is characterized by a fine and substantially uniform grain size distribution, preferably finer than ASTM 7 and more preferably in a range of about ASTM 8 to 10.

BACKGROUND OF THE INVENTION

The present invention generally relates to methods for processingnickel-base superalloys. More particularly, this invention relates to amethod of forging an article from a nickel-base superalloy, in whichincreased local strain rates in combination with increased carboncontent promote a more controlled grain growth during supersolvus heattreatment, such that the article is characterized by a microstructurewith a finer uniform grain size.

Gamma prime (γ′) precipitation-strengthened nickel-base superalloyscontain chromium, tungsten, molybdenum, rhenium and/or cobalt asprincipal elements that combine with nickel to form the gamma (γ)matrix, and contain aluminum, titanium, tantalum, niobium, and/orvanadium as principal elements that combine with nickel to form thedesirable gamma prime precipitate strengthening phase, principallyNi₃(Al, Ti). Gamma prime precipitation-strengthened nickel-basesuperalloys (hereinafter, gamma prime nickel-base superalloys) arewidely used for disks and other critical gas turbine engine componentsforged from billets produced by powder metallurgy (P/M), conventionalcast and wrought processing, and spraycast or nucleated casting formingtechniques. Gamma prime nickel-base superalloys formed by powdermetallurgy are particularly capable of providing a good balance ofcreep, tensile, and fatigue crack growth properties to meet theperformance requirements of certain gas turbine engine components suchas turbine disks. In a typical powder metallurgy process, a powder ofthe desired superalloy undergoes consolidation, such as by hot isostaticpressing (HIP) and/or extrusion consolidation. The resulting billet isthen isothermally forged at temperatures slightly below the gamma primesolvus temperature of the alloy to approach superplastic formingconditions, which allows the filling of the die cavity through theaccumulation of high geometric strains without the accumulation ofsignificant metallurgical strains. These processing steps are designedto retain the fine grain size originally within the billet (for example,ASTM 10 to 13 or finer), achieve high plasticity to fill near-net-shapeforging dies, avoid fracture during forging, and maintain relatively lowforging and die stresses. (Reference throughout to ASTM grain sizes isin accordance with the scale established in ASTM Standard E 112.) Inorder to improve fatigue crack growth resistance and mechanicalproperties at elevated temperatures, these alloys are then heat treatedabove their gamma prime solvus temperature (generally referred to assupersolvus heat treatment), to cause significant, uniform coarsening ofthe grains.

Forged gas turbine engine components often contain grains with sizes ofabout ASTM 9 and coarser, such as ASTM 2 to 9, though a much tighterrange is typically preferred, such as grain sizes within a limited rangeof 2 to 3 ASTM units. Such a limited range can be considered uniform,which as used herein refers to grain size and growth characterized bythe substantial absence of non-uniform critical grain growth. As usedherein, critical grain growth (CGG) refers to localized excessive graingrowth in an alloy that results in the formation of grains outsidetypical uniform grain size distributions whose size sufficiently exceedsthe average grain size in the alloy (such as regions as coarse as ASTM00 in a field of ASTM 6-10) to negatively affect the low cycle fatigue(LCF) properties of an article formed from the alloy, manifested byearly preferential crack nucleation in the CGG regions. Critical graingrowth can also have a negative impact on other mechanical properties,such as tensile strength. Critical grain growth occurs duringsupersolvus heat treatment following hot forging operations in which awide range of local strains and strain rates are introduced into thematerial. Though not wishing to be held to any particular theory,critical grain growth is believed to be driven by excessive storedenergy within the worked article, and may involve individual grains,multiple individual grains within a small region, or large areas ofadjacent grains. The grain diameters of the effected grains are oftensubstantially coarser than the desired grain size. Disks and othercritical gas turbine engine components forged from billets produced bypowder metallurgy and extrusion consolidation have appeared to exhibit alesser propensity for critical grain growth than if forged from billetsproduced by conventional cast and wrought processing or spraycastforming techniques, but in any event are susceptible to critical graingrowth during supersolvus heat treatment.

Commonly-assigned U.S. Pat. No. 4,957,567 to Krueger et al. teaches aprocess for eliminating critical (abnormal) grain growth in fine graingamma-prime nickel-base superalloy components by controlling thelocalized strain rates experienced during the hot forging operation.Strain rate is defined as the instantaneous rate of change of geometricstrain with time. Krueger et al. teach that local strain rates mustgenerally remain below a critical value, {dot over (ε)}_(c), in order toavoid detrimental critical grain growth during subsequent supersolvusheat treatment. According to Krueger et al., the maximum strain rate iscomposition, microstructure, and temperature dependent, and can bedetermined for a given superalloy by deforming test samples undervarious strain rate conditions, followed by a suitable supersolvus heattreatment. The maximum (critical) strain rate is then defined as thestrain rate that, if exceeded during deformation and working of asuperalloy and accompanied by a sufficient amount of total strain, willresult in critical grain growth after supersolvus heat treatment.

Another processing limitation identified by Krueger et al. as avoidingcritical grain growth in a nickel-base superalloy having a gamma primecontent of, for example, 30-46 volume percent and higher, is to ensuresuperplastic deformation of the billet during forging. For this purpose,the billet is processed to have a fine grain microstructure thatachieves a minimum strain rate sensitivity (m) of about 0.3 or greaterfor the superalloy within the forging temperature range. As known in theart, the ability of a fine grain billet to deform superplastically isdependent on strain rate sensitivity, and superplastic materials exhibita low flow stress as represented by the following equation:

σ=K{dot over (ε)} ^(m)

where σ is the flow stress, K is a constant, {dot over (ε)} is thestrain rate, and m is the strain rate sensitivity, with higher values ofm corresponding to greater superplasticity.

Further improvements in the control of final grain size have beenachieved with the teachings of commonly-assigned U.S. Pat. No. 5,529,643to Yoon et al. and U.S. Pat. No. 5,584,947 to Raymond et al. In additionto the requirement for superplasticity during forging (in other words,maintaining a high m value), Raymond et al. teach the importance of amaximum strain rate in combination with chemistry control, particularlythe carbon and/or yttrium content of the alloy to achieve grain boundarypinning in alloys having a gamma prime content of up to 65 volumepercent. In a particular example, Raymond et al. cites an upper limitstrain rate of below about 0.032 per second (s⁻¹) for a gamma primenickel-base superalloy identified as Alloy D and commercially known asRené 88DT (R88DT; U.S. Pat. No. 4,957,567). In addition to maintaining ahigh m value, Yoon et al. also identifies a maximum strain rate of notmore than about 0.032 s⁻¹, particularly in reference to forging an alloyidentified in Yoon et al. as Alloy A, which again is R88DT. Yoon et al.further place an upper limit on the maximum strain rate gradient duringforging, and requires extended annealing of the forging at a subsolvustemperature to remove stored strain energy prior to performing asupersolvus heat treatment. Finally, Yoon et al. achieve optimumsuperplasticity by forming the billet to have a grain size of finer thanabout ASTM 12, and maintaining the billet microstructure to achieve aminimum strain rate sensitivity of about m=0.3 within the forgingtemperature range.

While the teachings of Krueger et al., Yoon et al., and Raymond et al.have been largely effective in controlling critical grain growth,implementation of their teachings has generally required the use of veryslow ram speed control of the forging press head (generally with asimple linear decay versus stroke control scheme), coupled by simulativemodeling to translate the press head deformation rate into actual metalstrain rate as a function of temperature, constitutive property data forthe forging stock, die shape, and die or mult lubrication.

In addition to the absence of critical grain growth, mechanicalproperties of components forged from fine grain nickel-base superalloysfurther benefit from improved control of the grain size distribution toachieve a distribution and average grain size that are, respectively, asnarrow and fine as possible. Such a capability is particularlybeneficial for high temperature, high gamma prime content (e.g., about30 volume percent and above) superalloys, such as R88DT, for which adesired uniform grain size is generally not coarser than ASTM 6 for gasturbine disks. Though prior forging practices of the type describedabove have achieved grain sizes in a range of ASTM 5 to 8, less thanoptimal mechanical properties can still result. For example, low cyclefatigue life is known to decrease with coarser average grain sizes, evenif uniform. The impact of average grain size on low cycle fatigueproperties of supersolvus heat treated P/M superalloys is most apparentat low to intermediate temperatures, such as in a range of about 400° F.to about 750° F. (about 200° C. to about 400° C.). While the overalltemperature capability and balance of properties that P/M alloys offerare very attractive and relied on for the most advanced current engineapplications, even more benefit from these alloys could be obtained iftheir low cycle fatigue properties at low to intermediate temperaturescould be improved.

In view of the above, it would be desirable if more grain sizerefinement (finer average grain size) could be achieved in gas turbineengine forgings, along with the avoidance of localized critical graingrowth. It would be particularly desirable if both of these goals couldbe achieved using a process whose parameters are within practical limitsof manufacturability and cost constraints.

BRIEF SUMMARY OF THE INVENTION

The present invention provides a method of forming components from gammaprime nickel-base superalloys. The method entails formulating such asuperalloy to have a sufficiently high carbon content and forging thesuperalloy at sufficiently high local strain rates so that, following asupersolvus heat treatment, the component is characterized by a fine andsubstantially uniform grain size distribution, preferably an averagegrain size finer than ASTM 7 and more preferably in an average range ofabout ASTM 8 to 10. The present invention is further capable of avoidingcritical grain growth that would produce individual grains or smallregions of grains having grain sizes of more than five and preferablythree ASTM units coarser than the average grain size in the component,or large regions that are uniform in grain size but with a grain sizecoarser than a desired grain size range of about two ASTM units.

The method includes formulating the superalloy to have a compositionsuitable for producing forged polycrystalline articles subjected to hightemperatures and dynamic loads, a notable example of which is a turbinedisk of a gas turbine engine. An exemplary example of such an alloy isthe aforementioned gamma prime precipitation-strengthened nickel-basesuperalloy R88DT, though it is foreseeable that the teachings of thisinvention could be extended to other gamma prime nickel-base superalloysapproximating the mechanical properties of R88DT critical to a turbinedisk, such as low cycle fatigue life. In contrast to R88DT, thesuperalloy employed in the method of this invention is formulated tohave a carbon content of at least 0.045 weight percent, and preferablyin excess of 0.060 weight percent, which is the conventional upper limitfor carbon in R88DT. A billet is formed of the superalloy and worked ata temperature below the gamma prime solvus temperature of the superalloyso as to form a worked article. In particular, the billet is workedwhile maintaining strain rates as high as possible to control averagegrain size, but below an upper strain rate limit to avoid critical graingrowth. According to one aspect of the invention, the billet is notrequired to be worked superplastically, i.e., can have a strain ratesensitivity (m) of less than 0.3 at the working (e.g., forging)temperature. In fact, it is preferred to work the billetnon-superplastically to achieve the finest grain sizes. The workedarticle is then heat treated at a temperature above the gamma primesolvus temperature of the superalloy for a duration sufficient touniformly coarsen the grains of the worked article, after which theworked article is cooled at a rate sufficient to reprecipitate gammaprime within the worked article. The cooled worked article has anaverage grain size of not coarser than ASTM 6 and preferably not coarserthan 7 ASTM, for example, in a range of about ASTM 8 to 10.

A significant advantage of this invention is that, in addition toavoiding critical grain growth and avoiding the necessity to forgesuperplastically, the higher strain rate limit of the process window forworking the billet has been shown to achieve significant control of theaverage grain size in the component and achieve a uniform grain sizedistribution within a desired narrower range that is significantly finerthan previously possible. In this manner, mechanical properties of thecomponent, including low cycle fatigue and tensile strength, can beimproved. Though not wishing to be held to any particular theory, it isbelieved that formulating a superalloy such as R88DT to contain a carbonlevel above its conventional upper limit (0.060 weight percent) allowsthe use of strain rates beyond the upper strain rate limit of 0.010 persecond (s⁻¹) typically associated with R88DT, and even above the upperstrain rate limit of 0.032 s⁻¹ previously permitted by Yoon et al. andRaymond et al. for R88DT, resulting in components capable of exhibitinga more refined average grain size and substantially free of criticalgrain growth, which together improve the low cycle fatigue life of thecomponent. Low cycle fatigue life is particularly improved within atemperature range of about 400° F. to about 750° F. (about 200° C. toabout 400° C.) relative to R88DT with a conventional carbon content ofup to 0.060 weight percent.

Improvements in low cycle fatigue life are believed possible, with thefurther benefit of higher temperature properties achieved with powdermetallurgy alloys such as R88DT. Other benefits of the finer averagegrain size achieved with this invention include improved sonicinspection capability due to lower sonic noise, and improved yieldbehavior in service due to improved yield strength with finer grainsize.

Other objects and advantages of this invention will be betterappreciated from the following detailed description.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic graph representing strain rate versus temperatureand resulting grain size distribution for conventional R88DT specimensand R88DT specimens modified to contain carbon levels in excess of 0.060weight percent.

FIGS. 2 and 3 are graphs plotting average and ALA grain size,respectively, for R88DT specimens modified to contain either 0.066 or0.070 weight percent carbon and processed with relatively fast (at least0.01 s⁻¹) and slow (less than 0.01 s⁻¹) strain rate regimes.

FIG. 4 is a graph plotting strain rate sensitivity value and averagegrain size versus strain rate for data of FIG. 2 corresponding tospecimens forged at 1875° F.

FIG. 5 is a graph plotting the average and ALA grain size versus forgingtemperature data in FIGS. 2 and 3 for the R88DT specimens containing0.070 weight percent carbon.

FIGS. 6 and 7 are normal probability plots of the average and ALA grainsize, respectively, for specimens formulated and forged in accordancewith current practices and the present invention.

FIG. 8 is a bar graph comparing the sonicability of the forged specimensof FIGS. 6 and 7.

FIGS. 9 and 10 are normal probability plots of the ultimate tensilestrength and yield strength, size, respectively, of the forged specimensof FIGS. 6 and 7.

DETAILED DESCRIPTION OF THE INVENTION

The present invention is particularly directed to components formed byforging gamma prime precipitation-strengthened nickel-base superalloys.A particular example is high pressure turbine disks of gas turbineengines, which are typically formed by isothermally forging afine-grained billet at temperatures at or near the recrystallizationtemperature of the alloy but less than the gamma prime solvustemperature of the alloy, and under superplastic forming conditions toenable filling of the forging die cavity through the accumulation ofhigh geometric strains without the accumulation of significantmetallurgical strains. After forging, a supersolvus heat treatment isperformed, during which grain growth occurs. In the past, such asupersolvus heat treatment has typically yielded an acceptable but notwholly optimal average grain size range of about ASTM 2 to 9. Inaccordance with commonly-assigned U.S. Pat. No. 4,957,567 to Krueger etal., U.S. Pat. No. 5,529,643 to Yoon et al., and U.S. Pat. No. 5,584,947to Raymond et al., whose teachings regarding strain rates, strain rategradients, and superplasticity are incorporated herein by reference,placing an upper limit on the strain rate (critical strain rate), anupper limit on the strain rate gradient (critical strain rate gradient),and a strain rate sensitivity (m) of at least about 0.3 during forgingavoids critical grain growth during supersolvus heat treatment. However,even with the benefits of Krueger et al., Yoon et al., and Raymond etal., grain sizes in forged gamma prime precipitation-strengthenednickel-base superalloys have typically been limited to a range of aboutASTM 5 to 8.

The present invention identifies processing parameters by which a fineraverage grain size and more desirable grain size distribution can beachieved in a gamma prime precipitation-strengthened nickel-basesuperalloy, in addition to avoidance of critical grain growth. Accordingto one aspect of the invention, a finer and more controllable averagegrain size can be achieved by increasing the strain rate during forging,resulting in a strain rate window having an upper limit that is higherthan conventionally believed possible without inducing critical graingrowth. The upper limit of the strain rate window corresponds to themaximum strain rate at which critical grain growth can be avoided.According to a second aspect of the invention, higher strain rates andnon-superplastic deformation can be employed without causing criticalgrain growth by modifying an alloy to have a relatively high carboncontent.

The above-noted aspects of the invention will be discussed in referenceto processing of a high-pressure turbine disk for a gas turbine engine.However, those skilled in the art will appreciate that the teachings andbenefits of this invention are applicable to numerous other componentsthat undergo forging and in which a fine controlled grain sizedistribution is desired.

In the production of a high pressure turbine disk from a gamma primenickel-base superalloy, a billet is typically formed by powdermetallurgy (P/M), a cast and wrought processing, or a spraycast ornucleated casting type technique. Such processes are carried out toyield a billet with a fine grain size, typically about ASTM 10 or finer,to achieve low flow stresses during forging. As previously noted, theability of a fine grain billet to deform superplastically is dependenton strain rate sensitivity (m). Whether formed by powder metallurgy,spraycast forming, cast and wrought, or another suitable method, priorart billets for high pressure turbine disks have been formed underconditions, including a specified temperature range, to produce thedesired fine grain size and also maintain a minimum strain ratesensitivity (m) of about 0.3 or greater within the forging temperaturerange. Alternatively, to control the strain rate sensitivity, it hasbeen conventional practice to control the forging process to besuperplastic by forging in a regime of strain rate and temperature whereflow stress is constant for any strain (negligible strain hardening orstrain softening). As will be discussed in more detail below, with thepresent invention it is believed that strain rates higher thanpreviously thought possible can be employed to produce suitable forgingswithout the forging process being fully superplastic, i.e., at strainrate sensitivity values of less than about 0.3.

In a preferred embodiment utilizing a powder metallurgy process, thebillet can be formed by consolidating a superalloy powder, such as byhot isostatic pressing (HIP) or extrusion consolidation, the latter ofwhich preferably uses a sufficiently low ram speed to prevent adiabaticheating and limited only by equipment tonnage limitations and excessivechilling. As known in the art, consolidation preferably yields a fullydense, fine-grain billet preferably having at least about 98%theoretical density. Prior to working the billet, a high temperaturesoak is typically performed in a manner that prevents excessivecoarsening of the overall grain size that would excessively andundesirably increase flow stresses. In the practice of the presentinvention, a suitable soak has been achieved by simply preheating andholding the billet at its forging temperature for up to about fivehours, though longer holding periods are also envisioned.

The billet is then hot worked (e.g., forged) to form a component havinga desired geometry, followed by a supersolvus (solution) heat treatment.As taught by Yoon et al., under certain conditions an extended subsolvusannealing process or a low heating rate to the supersolvus heattreatment temperature may be desired to dissipate stored strain energywithin the article and equilibrate the temperature of the component.Dissipation of stored strain energy can serve to reduce nonuniformnucleation tendencies of the superalloy, such that the tendency forcritical grain growth in the component is also reduced. However, in thepresent invention such an extended subsolvus anneal step appears to beunnecessary. Instead, merely preheating the worked billet (forging) towithin about 50° F. to about 75° F. (about 30° C. to about 40° C.) or soof the prior forging temperature is sufficient without any extended soaktime.

The supersolvus heat treatment is then performed at a temperature abovethe gamma prime solvus temperature (but below the incipient meltingtemperature) of the superalloy, to recrystallize the worked grainstructure and dissolve (solution) the gamma prime precipitates in thesuperalloy. To accommodate furnace variability, a suitable supersolvustemperature is typically about 30° F. to 70° F. (about 15° C. to 40° C.)above the gamma prime solvus temperature of an alloy, though anytemperature above the solvus temperature (but below the incipientmelting temperature) is generally acceptable. Following the supersolvusheat treatment, the component is cooled at an appropriate rate tore-precipitate gamma prime within the gamma matrix or at grainboundaries, so as to achieve the particular mechanical propertiesdesired. An example of a suitable cooling step includes controlled aircooling or controlled air cooling for a brief period followed byquenching in oil or another suitable medium. The component may also beaged using known techniques with a short stress relief cycle at atemperature above the aging temperature of the alloy if desirable toreduce residual stresses.

As noted above and well known in the art, in addition to grainrecrystallization and solutioning gamma prime precipitates, heating thesuperalloy above its gamma prime solvus temperature causes grain growth(coarsening), typically resulting in grain sizes coarser than theoriginal billet grain size, for example, coarser than about ASTM 9, suchas in a range of about ASTM 2 to 9. To achieve mechanical propertiesdesired for a gas turbine disk, uniform average grain sizes within arange of about two or three ASTM units are typically desired. Regions ofthe component with grain sizes in excess of about two to three ASTMunits coarser than the desired grain size range are undesirable in thatthe presence of such grains can significantly reduce the low cyclefatigue resistance of the component and have a negative impact on othermechanical properties of the component, such as tensile and fatiguestrength. For example, a component having a grain size range of aboutASTM 5 to 8 is preferably free of isolated grains and small regions ofgrains coarser than ASTM 3 (though widely scattered grains slightlycoarser may be tolerable), and free of significant regions coarser thanabout ASTM 6. As noted above, excessively large grains caused bycritical grain growth can be avoided during working of the billet bymaintaining strain rates below a critical (maximum) strain rate for thesuperalloy in accordance with Krueger et al. However, mechanicalproperties would be further promoted by improving the grain sizedistribution and achieving a finer average grain size, for example, in arange of about ASTM 7 to 9, more preferably about 8 to 10.

According to the present invention, improved grain size distribution andfiner average grain size can be achieved by increasing the minimumstrain rate of the strain rate window. Furthermore, the maximum strainrate can be increased to values previously associated with causingcritical grain growth in a given superalloy, but without inducingcritical grain growth, by increasing the carbon content of thesuperalloy above its conventional upper limit. In part, the effect ofthe increased carbon content is believed to be an increased pinningforce that inhibits abnormal grain growth. Generally, finely dispersedcarbides restrict grain boundary motion during supersolvus heattreatment, such that the grains are not permitted to grow excessivelyand/or randomly to the extent that critical grain growth occurs. Fromthe investigations reported below, in addition to a more rapid forgingprocess and improved properties, other benefits appear to be the abilityto perform the forging operation at relatively low temperatures andunder non-superplastic conditions (m<0.3).

According to Krueger et al., the critical strain rate of a gamma primenickel-base superalloy is composition, microstructure, and temperaturedependent, and can be determined for a given superalloy by deformingtest samples under various strain rate conditions, and then performingsuitable supersolvus heat treatments. The critical strain rate is thendefined as the strain rate that, if exceeded during deformation andworking of a superalloy and accompanied by a sufficient amount of totalstrain, will result in critical grain growth after supersolvus heattreatment. According to the present invention, in which higher strainrates are identified as being capable of achieving a more controlled andfiner average grain size after supersolvus heat treatment, strain ratesbelow a minimum strain rate result in an average grain size that may becoarser than desired for optimal properties. As with the maximum strainrate identified by Krueger et al., the precise value for the minimumstrain rate parameter of this invention appears to vary depending on thecomposition and microstructure of the superalloy in question. Strainrates for regions within large components can be predicted analyticallyby performing experiments on small laboratory specimens, and then usingmodeling techniques to predict local deformation behavior within thecomponents.

The ability to improve grain size distribution and achieve finer averagegrain size by increasing the strain rate during forging is representedin FIG. 1, which is a graph representing strain rate versus temperatureand resulting average grain sizes observed with forged specimens of theconventional R88DT alloy (“Prior art forgings”) and forged specimens ofalloys based on the R88DT but whose compositions were modified tocontain elevated carbon levels in accordance with this invention (“Finegrain forgings”). Specifically, the conventional and modified R88DTalloys differ in their carbon contents, with specimens of theconventional R88DT alloy containing about 0.045 to 0.060 weight percentcarbon and specimens of the modified R88DT alloys containing greaterthan 0.060 weight percent carbon, preferably at least 0.065% to about0.085% carbon, and possibly higher. The bars in FIG. 1 represent a shiftin acceptable ranges for strain rates and forging temperaturesidentified through investigations leading to the present invention. Theupper extent of each bar represents the upper strain rate limit at whichcritical grain growth can be avoided in the specimens represented bythat bar. From FIG. 1, it can be appreciated that the modified R88DTspecimens are able to be forged at much higher rates than theconventional R88DT specimens without critical grain growth (about 0.1s⁻¹ maximum versus about 0.010 s⁻¹ maximum). FIG. 1 also represents thatsignificantly finer grains were obtained with the modified R88DTspecimens (about ASTM 7 to about ASTM 9) as compared to the conventionalR88DT specimens (about ASTM 6 to about ASTM 8).

The parameters and effects represented in FIG. 1 were determined througha series of investigations. In preliminary investigations, therelationship between final grain size and strain rate, including maximum(critical) strain rate, was evidenced from testing performed on subscaleright circular cylinder (RCC) and double cone (DC) specimens. Allspecimens were formed from compositions based on the superalloy R88DT,which is disclosed in commonly-assigned U.S. Pat. No. 4,957,567 toKrueger et al. as having a composition of, by weight, about 15.0-17.0%chromium, about 12.0-14.0% cobalt, about 3.5-4.5% molybdenum, about3.5-4.5% tungsten, about 1.5-2.5% aluminum, about 3.2-4.2% titanium,about 0.5.0-1.0% niobium, about 0.010-0.060% carbon, about 0.010-0.060%zirconium, about 0.010-0.040% boron, about 0.0-0.3% hafnium, about0.0-0.01 vanadium, and about 0.0-0.01 yttrium, the balance nickel andincidental impurities. The gamma prime solvus temperature of R88DT isestimated to be about 1950-2150° F. (about 1065-1180° C.), typically2025-2050° F. (about 1105-1120° C.), for about 40 volume percent gammaprime. The actual chemistries of the specimens are summarized in thetable below.

“0.066% C Specimens” “0.070% C Specimens” Chromium 15.92 weight percent15.86 weight percent Cobalt 12.78 12.90 Molybdenum 3.92 3.93 Tungsten3.98 3.90 Aluminum 2.13 2.16 Titanium 3.74 3.68 Niobium 0.70 0.67Zirconium 0.40 0.039 Boron 0.014 0.014 Hafnium 0.00015 0.000039 Vanadium0.0026 0.0026 Yttrium <0.0005 <0.0005 Iron 0.04 0.03 Carbon 0.066 0.070Nickel Balance (+ impurities) Balance (+ impurities)From the above, it can be seen that the investigation evaluated twogroups of experimental specimens alloyed to contain either about 0.066%or about 0.070% carbon.

FIGS. 2 and 3 plot, respectively, average ASTM grain size (ASTM StandardE 112) and ALA grain size (ASTM Standard E 930) versus strain rates forboth groups of RCC specimens. All specimens were forged at a temperatureof about 1850° F., 1875° F., or 1900° F. (about 1010° C., about 1025°C., or about 1040° C.), and at a strain rate within a range of about0.00032 to about 1 sec⁻¹. Nominal strain levels were about 0.7%. Forgingtemperatures were selected on the basis of prior investigations withsimilar RCC and DC specimens of the conventional R88DT alloy (the basisfor the conventional R88DT data represented in FIG. 1). Though priorwork reported by Huron, “Control of Grain Size via Forging Strain RateLimits for R88DT,” Superalloys 2000, Sep. 17-21, 2000 (a publication ofTMS), has shown that a strain rate of about 0.010 sec⁻¹ is an upperstrain rate limit for specimens of the conventional R88DT alloy in orderto avoid critical grain growth, critical grain growth was not observedin any of the experimental specimens at any of the tested strain rates,even at strain rates of about 1 sec⁻¹ (0.066% C specimens forged at1875° F.), which is a factor of one thousand over what was previouslythought possible. In contrast, the data corresponding to the currentlower strain rate practice of rates less than 0.010 sec⁻¹ evidence theentire grain size distributions, for both average and ALA, are shiftedto coarser grain sizes, regardless of carbon content, attesting to theimpact of the strain rate effect.

From the above results, it was concluded that the elevated carboncontents of the experimental specimens enabled forging at strain ratessignificantly higher than was previously possible with the conventionalR88DT alloy, and even higher than thought possible for the high carbonalloys taught by Raymond et al., without encountering critical graingrowth. It was further concluded that the higher strain rates resultedin a finer average grain size than was previously possible with theconventional R88DT alloy. The average grain sizes achieved within thespecimens reported in FIG. 2 are sufficiently finer to have a beneficialand significant effect on low cycle fatigue life and other mechanicalproperties such as ultimate tensile strength (UTS), as well asprocessing considerations such as reduced noise during sonicinspections.

The data plot of FIG. 3 evidences ALA grain size also shows improvementover current lower strain rate practice at rates less than 0.010 sec⁻¹.The data show that, even at higher carbon levels, the entire grain sizedistributions (both average and ALA) are shifted to coarser grain sizes,again attesting to the impact of the strain rate effect.

While the data plotted in FIGS. 2 and 3 are for specimens containingeither 0.066 or 0.070 weight percent carbon, the trends evidenced in thegraphs indicate that similar benefits and even further improvements arepossible with higher carbon contents. Based on these results, a suitablerange for the carbon content is believed to be about 0.065 to about0.085 weight percent. Higher carbon contents are also believed to bepossible, such as about 0.10 weight percent, with the upper limit beinggenerally limited only by the potential detrimental impact of excessivecarbon on other properties of the superalloy.

As noted above, with the present invention it is believed that higherstrain rates can be employed to produce suitable forgings without theforging process being fully superplastic, in other words, forging can beperformed with strain rate sensitivity values of less than about 0.3.This aspect of the invention is evident from FIG. 4, which is a graphplotting the strain rate sensitivity value (m) and average grain sizeversus strain rate for the data of FIG. 2 corresponding to the 0.066% Cspecimens forged at 1875° F. FIG. 4 evidences that finer average grainsizes were achieved with strain rate sensitivity values of less thanabout 0.3, and therefore well below the minimum strain rate sensitivitytaught by Yoon et al. and Raymond et al.

FIG. 5 is a graph plotting average and ALA grain size versus forgingtemperature for the 0.070% C specimens. FIG. 5 suggests that a trendalso exists for finer and more uniform average grain sizes withdecreasing forging temperature in the 0.070% C specimens. This trend wasalso found for the 0.066% C specimens, evidencing a broad processingwindow for high carbon contents and the potential benefits of lowerforging temperatures.

In another investigation, high pressure turbine disks were forged andanalyzed to further assess the above-described findings regarding theability to obtain finer average grain size by increasing strain ratesduring forging. Three of the disks were formed from R88DT modified tohave a carbon content of either about 0.066 or about 0.070 weightpercent, and were produced by powder metallurgy, extrusionconsolidation, forging, and supersolvus heat treatment at about 2080° F.(about 1140° C.). As understood by those skilled in the forging art,forging processes can be designed using simulation models to produce dieshapes and achieve a forging press operation that controls the localstrain and strain rate history of regions of a forging within desiredparameters. Furthermore, because the forged disks of this investigationand the forged specimens of the previous investigations have an array oflocal strains and strain rates for a given macro overall strain ratebased on the upset ratio of the workpiece, the forge rates for the disksand previous specimens can be compared on the basis of a maximum strainrate. Using this approach, the 0.066% C and 0.070% C disks evaluated inthis investigation were forged using nominally isothermal processesdesigned to achieve maximum strain rates of about 0.032 sec⁻¹. Theforging steps for the 0.066% C and 0.070% C disks were controlled on alocal limit basis so that all regions of the forgings were at or belowthe 0.032 sec⁻¹ upper limit.

The consistency with which finer grain sizes can be obtained with higherstrain rates is evidenced in FIGS. 6 and 7, which are normal probabilityplots of the average and ALA grain size, respectively, of the threeforgings (Forgings #2, #3, and #4) produced to have one of theabove-noted 0.066% C, and 0.070% C compositions. Twenty measurementswere obtained for each forging arrayed uniformly about the forging crosssection. The data for the 0.066% C and 0.070% C forgings are plottedalong with data obtained from a fourth disk of the same geometry, butformed from a conventional R88DT composition including a conventionalcarbon content of about 0.052 weight percent, and processed to theconventional strain rate limit of less than 0.010 sec⁻¹ (Forging #1).Even with the complexity of geometric factors in a contoured forging,the improved practice of the invention demonstrates a finer mean averagegrain size by about 2 ASTM grain size numbers and an improvement in themean ALA grain size by about 1 ASTM grain size number.

Based on these results, to achieve the benefit in complex contouredforgings where geometric factors drive local strain rate variations, itis believed that maximum strain rates of about 0.032 sec⁻¹ and above,which correspond to strain rate sensitivity values of less than 0.3,should be employed to achieve the refined grain size throughout suchforgings. As previously noted, it is very unlikely in a complex forgingthat a target maximum strain rate will be uniformly achieved in allareas of the forging, and variations in strain rate can be such thatsetting an absolute minimum strain rate in the forging is not practical.On the other hand, maximum strain rates capable of avoiding criticalgrain growth while achieving a finer grain size and distribution inaccordance with the invention will inherently fall over a range.Therefore, a maximum strain rate can be set as a target within a rangeof suitable maximum strain rates for a given forging, in which an upperlimit for the range is necessary to avoid critical grain growth and alower limit of the range is necessary to avoid or minimize low strainareas that may not achieve sufficiently high strain rate work to obtainthe desired fine grain size and distribution sought by the invention.Alternatively or in addition, the forging shape may be defined so thatthe high strain rate non-superplastically deformed regions are locatedin specific areas advantageous to the part operation and life.

The average grain sizes achieved within the specimens reported in FIGS.6 and 7 were concluded to be sufficiently finer to have a beneficial andsignificant effect on low cycle fatigue life and other mechanicalproperties such as ultimate tensile strength (UTS), as well asprocessing considerations such as reduced noise during sonicinspections. Such benefits are evident in FIGS. 8 through 10. FIG. 8 isa bar graph evidencing the improved sonicability of the 0.066% C and0.070% C forgings of FIGS. 6 and 7 as compared to the conventional0.052% C forging of FIGS. 6 and 7. The data show a reduction of sonicnoise levels of about 40%, indicating improved inspectability comparedto the conventional current processing and chemistry.

Finally, FIGS. 9 and 10 are normal probability plots comparing theultimate tensile strength and yield strength, respectively, of the0.066% C and 0.070% C forgings of FIGS. 6 and 7 as compared to theconventional 0.052% C forging of FIGS. 6 and 7. Both measures of tensilecapability (ultimate tensile strength and 0.2% yield strength) show asignificant improvement of the mean value of about 9 to about 10 ksi(about 62 to about 69 MPa), when using methods and compositions of thisinvention as compared to currently existing methods and compositions.

In view of the above, the method of this invention makes possible theproduction of components from R88DT and similar gamma prime nickel-basesuperalloys that consistently exhibit a finer average grain size. Whilethe benefits of the invention were described in reference to the R88DTsuperalloy processed from powder metal starting materials, othermaterials could be used including spraycast materials, cast and wroughtmaterials, etc. Furthermore, gamma prime nickel-base superalloys havingcompositions sufficiently approximating that of R88DT to have similarmechanical properties as R88DT, such as low cycle fatigue life, are alsobelieved to benefit from the processing and composition modifications ofthe present invention. An example of such an alloy is believed to beRené 104 (U.S. Pat. No. 6,521,175), with a nominal composition of, byweight, about 16.0-22.4% cobalt, about 6.6-14.3% chromium, about2.6-4.8% aluminum, about 2.4-4.6% titanium, about 1.4-3.5% tantalum,about 0.9-3.0% niobium, about 1.9-4.0% tungsten, about 1.9-3.9%molybdenum, about 0.0-2.5% rhenium, about 0.02-0.10% carbon, about0.02-0.10% boron, about 0.03-0.10% zirconium, the balance nickel andincidental impurities. Another notable example is NF3 (U.S. Pat. No.6,521,175), with a nominal composition of about 16.0-20.0 percentcobalt, about 8.5-12.5 percent chromium, about 1.5-3.5 percent tantalum,about 2.0-4.0 percent tungsten, about 1.9-3.9 percent molybdenum, about0.04-0.06 percent zirconium, about 1.0-3.0 percent niobium, about2.6-4.6 percent titanium, about 2.6-4.6 percent aluminum, about0.02-0.04 percent carbon, about 0.02-0.04 percent boron, the balancenickel and incidental impurities.

While the invention has been described in terms of particular processingparameters and compositions, the scope of the invention is not solimited. Instead, modifications could be adopted by one skilled in theart, such as by substituting other gamma prime precipitationstrengthened nickel-base superalloys with higher or lower gamma primecontents, or by modifying the preferred method by substituting otherprocessing steps or including additional processing steps. Accordingly,the scope of the invention is to be limited only by the followingclaims.

1. A method of forming an article from a gamma primeprecipitation-strengthened nickel-base superalloy having a gamma primesolvus temperature, the method comprising the steps of: formulating thegamma prime precipitation-strengthened nickel-base superalloy to containgreater than 0.060 weight percent carbon; forming a billet of thesuperalloy; working the billet at a temperature below the gamma primesolvus temperature of the superalloy so as to form a worked article,wherein the billet is worked to undergo non-superplastic deformation andto achieve a maximum strain rate that is below an upper strain ratelimit to avoid critical grain growth yet sufficiently high to controlaverage grain size, wherein the upper strain rate limit is greater than0.008 per second; heat treating the worked article at a temperatureabove the gamma prime solvus temperature of the superalloy for aduration sufficient to uniformly coarsen the grains of the workedarticle; and cooling the worked article at a rate sufficient toreprecipitate gamma prime within the worked article, wherein the workedarticle has an average grain size of not coarser than ASTM 7 and issubstantially free of grains in excess of three ASTM units coarser thanthe average grain size.
 2. The method according to claim 1, wherein theforming step comprises a process chosen from the group consisting ofpowder metallurgy, cast and wrought, and spraycast forming techniques.3. The method according to claim 1, wherein the forming step compriseshot isostatic pressing or extrusion consolidation of a powder of thesuperalloy to form the billet.
 4. The method according to claim 1,wherein the superalloy contains up to about 0.10% carbon.
 5. The methodaccording to claim 1, wherein the superalloy contains 0.065% to 0.085%carbon.
 6. The method according to claim 1, wherein the superalloycontains 0.066% to 0.070% carbon.
 7. The method according to claim 1,wherein the maximum strain rate is at least 0.010 per second.
 8. Themethod according to claim 1, wherein the maximum strain rate is at least0.032 per second.
 9. The method according to claim 1, wherein the upperstrain rate limit is greater than 0.100 per second.
 10. The methodaccording to claim 1, wherein the worked article has an average grainsize in a range of about ASTM 7 to
 10. 11. The method according to claim1, wherein the worked article has an average grain size of not coarserthan ASTM
 8. 12. The method according to claim 1, wherein the workedarticle has an average grain size of about ASTM 8 to about ASTM
 10. 13.The method according to claim 1, wherein the billet working step ischaracterized by a minimum strain rate sensitivity of less than m=0.3.14. The method according to claim 1, wherein the billet is worked sothat nominal strain within the billet is about 0.7.
 15. The methodaccording to claim 1, wherein the superalloy contains, by weight, about15.0-17.0% chromium, about 12.0-14.0% cobalt, about 3.5-4.5% molybdenum,about 3.5-4.5% tungsten, about 1.5-2.5% aluminum, about 3.2-4.2%titanium, about 0.5.0-1.0% niobium, at least 0.065-0.10% carbon, about0.010-0.060% zirconium, about 0.010-0.040% boron, about 0.0-0.3%hafnium, about 0.0-0.01 vanadium, and about 0.0-0.01 yttrium, thebalance essentially nickel and incidental impurities.
 16. The workedarticle formed by the method of claim 15, wherein the worked article isa turbine disk of a gas turbine engine, and after the cooling step theworked article has an average grain size of ASTM 8 to
 10. 17. The methodaccording to claim 1, wherein the superalloy contains, by weight, about16.0-22.4% cobalt, about 6.6-14.3% chromium, about 2.6-4.8% aluminum,about 2.4-4.6% titanium, about 1.4-3.5% tantalum, about 0.9-3.0%niobium, about 1.9-4.0% tungsten, about 1.9-3.9% molybdenum, about0.0-2.5% rhenium, about 0.02-0.10% carbon, about 0.02-0.10% boron, about0.03-0.10% zirconium, the balance essentially nickel and incidentalimpurities.
 18. The worked article formed by the method of claim 17,wherein the worked article is a turbine disk of a gas turbine engine,and after the cooling step the worked article has an average grain sizeof ASTM 8 to
 10. 19. The method according to claim 1, wherein thesuperalloy contains, by weight, about 16.0-20.0 percent cobalt, about8.5-12.5 percent chromium, about 1.5-3.5 percent tantalum, about 2.0-4.0percent tungsten, about 1.9-3.9 percent molybdenum, about 0.04-0.06percent zirconium, about 1.0-3.0 percent niobium, about 2.6-4.6 percenttitanium, about 2.6-4.6 percent aluminum, about 0.02-0.04 percentcarbon, about 0.02-0.04 percent boron, the balance essentially nickeland incidental impurities.
 20. The worked article formed by the methodof claim 19, wherein the worked article is a turbine disk of a gasturbine engine, and after the cooling step the worked article has anaverage grain size of ASTM 8 to
 10. 21. The worked article formed by themethod of claim 1, wherein the worked article is a turbine disk of a gasturbine engine, and after the cooling step the worked article has anaverage grain size of ASTM 8 to
 10. 22. A method of forming an articlefrom a gamma prime precipitation-strengthened nickel-base superalloyhaving a gamma prime solvus temperature, the method comprising the stepsof: formulating the superalloy to consist of, by weight, about15.0-17.0% chromium, about 12.0-14.0% cobalt, about 3.5-4.5% molybdenum,about 3.5-4.5% tungsten, about 1.5-2.5% aluminum, about 3.2-4.2%titanium, about 0.5.0-1.0% niobium, at least 0.065% to about 0.10%carbon, about 0.010-0.060% zirconium, about 0.010-0.040% boron, about0.0-0.3% hafnium, about 0.0-0.01 vanadium, and about 0.0-0.01 yttrium,the balance nickel and incidental impurities; forming a billet of thesuperalloy to have a fine grain size; working the billet at atemperature below the gamma prime solvus temperature of the superalloyso as to form a worked article, the billet working step beingcharacterized by a minimum strain rate sensitivity of less than m=0.3 soas not to be fully superplastic during the working step, the workingstep being performed to achieve a maximum strain rate that is below anupper strain rate limit to avoid critical grain growth yet sufficientlyhigh to control average grain size, wherein the maximum strain rate isat least 0.010 per second and the upper strain rate limit is greaterthan 0.1 per second; heat treating the worked article at a temperatureabove the gamma prime solvus temperature of the superalloy for aduration sufficient to uniformly coarsen the grains of the workedarticle; and cooling the worked article at a rate sufficient toreprecipitate gamma prime within the worked article, wherein the workedarticle has an average grain size of not coarser than ASTM 7 and issubstantially free of grains in excess of two ASTM units coarser thanthe average grain size.
 23. The method according to claim 22, whereinthe maximum strain rate is at least 0.032 per second.
 24. The methodaccording to claim 22, wherein the superalloy contains 0.065% to 0.085%carbon.
 25. The method according to claim 22, wherein the superalloycontains 0.066% to 0.070% carbon.
 26. The worked article formed by themethod of claim 22, wherein the worked article is a turbine disk of agas turbine engine, and after the cooling step the worked article has anaverage grain size of about ASTM 8 to 10.